Graphene is at the centre of nanotechnology research. In order to fully exploit its outstanding properties, a mass production method is necessary. Two main routes are possible: large-scale growth or large-scale exfoliation. Here, we demonstrate graphene dispersions with concentrations up to ~0.01 mg/ml by dispersion and exfoliation of graphite in organic solvents such as N-methylpyrrolidone. This occurs because the energy required to exfoliate graphene is balanced by the solvent-graphene interaction for solvents whose surface energy matches that of graphene. We confirm the presence of individual graphene sheets with yields of up to 12% by mass, using absorption spectroscopy, transmission electron microscopy and electron diffraction. The absence of defects or oxides is confirmed by X-ray photoelectron, infra-red and Raman spectroscopies. We can produce conductive, semi-transparent films and conductive composites. Solution processing of graphene opens up a whole range of potential large-scale applications from device or sensor fabrication to liquid phase chemistry. Hernandez et al 2Graphene is one of the most exciting nano-materials due to the cascade of unique physical properties that have recently been demonstrated. For example, due to the details of its electronic structure, charge carriers in graphene behave as massless Dirac fermions 1 . Furthermore, novel effects such as an ambipolar field effect 2 , room temperature quantum Hall effect 3 , breakdown of the Born-Oppenheimer approximation 4 are observed. However, as was the case in the early days of nanotube and nanowire research, graphene at present still suffers from one problem, critical for its mass-scale exploitation: it cannot yet be made with high yield. The standard procedure used to make graphene is micromechanical cleavage 5 . This yields the best samples to date, with mobilities up to 200,000 cm 2 /Vs. 6 However, single layers are a negligible fraction amongst large quantities of thin graphite flakes. Furthermore, it is difficult to see how to scale up this process to mass production. Alternatively, growth of graphene is also commonly achieved by annealing SiC substrates, but these samples are in fact composed of a multitude of domains, most of them sub-micrometer, and not spatially uniform in number, or in size over larger length scales 7 . A number of works have also reported graphene growth on metal substrates 8,9 , but this would require the sample transfer to insulating substrates in order to make useful devices, either via mechanical transfer or, via solution processing.Recently, a large number of papers have described the dispersion and exfoliation of graphene oxide (GO) [10][11][12][13] . This material consists of graphene-like sheets, chemically functionalised with compounds such as hydroxyls and epoxides, which stabilise the sheets in water 14 . However, this functionalisation results in considerable disruption of the electronic structure of the graphene. In fact GO is an insulator 15 rather than a semi-metal and is conceptually differen...

We have demonstrated a method to disperse and exfoliate graphite to give graphene suspended in water-surfactant solutions. Optical characterisation of these suspensions allowed the partial optimisation of the dispersion process. Transmission electron microscopy showed the dispersed phase to consist of small graphitic flakes. More than 40% of these flakes had <5 layers with ~3% of flakes consisting of monolayers. These flakes are stabilised against reaggregation by Coulomb repulsion due to the adsorbed surfactant. However, the larger flakes tend to sediment out over ~6 weeks, leaving only small flakes dispersed. It is possible to form thin films by vacuum filtration of these dispersions. Raman and IR spectroscopic analysis of these films suggests the flakes to be largely free of defects and oxides. The deposited films are reasonably conductive and are semi-transparent. Further improvements may result in the development of cheap transparent conductors.

We have characterized both the direct current conductivity and morphology of a wide range of films made from bundled nanotubes, produced by a selection of commercial suppliers. The conductivity increases with increasing nanotube graphitization but decreases with increasing film porosity P and mean bundle diameter ͗D͘. Computational studies show that the network conductivity is expected to scale linearly with the number density of interbundle junctions. A simple expression is derived to relate the junction number density to the porosity and mean bundle diameter. Plotting the experimental network conductivities versus the junction number density calculated from porosity and bundle diameter shows an approximate linear relationship. Such a linear relationship implies that the conductivity scales quadratically with the nanotube volume fraction, reminiscent of percolation theory. More importantly it shows the conductivity to scale with ͗D͘ −3 . Well-defined scaling with diameter and porosity allows the calculation of a specific conductivity expected for films with porosity of 50% and mean bundle diameter of 2 nm. This predicted specific conductivity scales well with the level of nanotube graphitization, reaching values as high as 1.5ϫ 10 7 S / m for well graphitized HiPCO single walled nanotubes.

We describe the fabrication of extremely high mass fraction polymer‐nanotube composites displaying very high conductivity. Contrary to general expectations, the conductivity displays percolation‐like scaling for all mass fractions from 22 % right up to 100 %. These are among the most conductive composites ever demonstrated.

We have prepared polyvinylalcohol-SWNT fibers with diameters from ∼1 to 15 μm by coagulation spinning. When normalized to nanotube volume fraction, V(f), both fiber modulus, Y, and strength, σ(B), scale strongly with fiber diameter, D: Y/V(f) ∝ D(-1.55) and σ(B)/V(f) ∝ D(-1.75). We show that much of this dependence is attributable to correlation between V(f) and D due to details of the spinning process: V(f) ∝ D(0.93). However, by carrying out Weibull failure analysis and measuring the orientation distribution of the nanotubes, we show that the rest of the diameter dependence is due to a combination of defect and orientation effects. For a given nanotube volume fraction, the fiber strength scales as σ(B) ∝ D(-0.29)D(-0.64), with the first and second terms representing the defect and orientation contributions, respectively. The orientation term is present and dominates for fibers of diameter between 4 and 50 μm. By preparing fibers with low diameter (1-2 μm), we have obtained mean mechanical properties as high as Y = 244 GPa and σ(B) = 2.9 GPa.

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We have prepared composites from a thermoplastic polyurethane reinforced with functionalized single walled nanotubes. Nanotubes with two types of functional groups were used: water-soluble tubes functionalized with polyethyleneglycol or poly(amino benzene sulfonic acid) and tetrahydrafuran-soluble tubes functionalized with octadecylamine. Composites prepared with water-or tetrahydrafuran-soluble tubes showed markedly different properties. Addition of water-soluble tubes tended to result in crystallization of the polyurethane soft segments, whereas addition of the tetrahydrafuran-soluble tubes promoted crystallization of the polyurethane hard segments. We interpret this as evidence of selective insertion of tubes in either hard or soft segments depending on the surface chemistry of the (functionalized) nanotube and the chemical structure of the segment. This interpretation is supported by differences in the mechanical properties of the composites. The waterbased composites tend to be stiffer and display higher plateau stress, consistent with reinforcement of the soft segments. However, the tetrahydrafuran cast composites tend to maintain their strength and ductility at higher nanotube loading levels, whereas the water-based composites become weak and brittle above ∼10 vol % nanotubes. This is consistent with the water-based nanotubes impeding the extension and motion of the soft segments, resulting in loss of ductility. In contrast, the tetrahydrafuran-soluble nanotubes become segregated in the hard segments and so do not negatively impact on the mechanical properties at high nanotube content. This controlled reinforcement has allowed us to prepare composites with modulus, plateau stress, strength, and ductility of up to 250 MPa, 8 MPa, 60 MPa and 750%, respectively, significantly better than neat polyurethane.

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